the stability of microstructure in the iron-carbon system during cyclic deformation

THE STABILITY OF MICROSTRUCTURE IN THE IRON-CARBON SYSTEM DURING CYCLIC DEFORMATION*

M. J. BROWNat and J. I). EMBURY?

The repose of various cttrbide dispersions to cyclic deformation btts been es&mined for quench sged low carbon iron using fatigue tests &t constant, strain amplitude. In addition, the response of quenched, deformed and aged specimens has been investigated to determine the role of carbides in stabilising the dislocation microstructure against fatigue softening. The differences in dislooation arrangement after cyclic deformation and unidirectionel deformation sre rationalised in terms of the dissipation of localised strain gradients on reversed deformation. The rates of fatigue softening are considered to reflect the magnitude of the retaining forces exerted OR the dislocation substruoture by the dispersion of the second phase particles,

LA STABILITZ DE ~ICRUST~~~RE DAh’S LE SYST&IE I?ER-CARBOXE df COURS DE D&FOR&ATIOhT CYCLIQUE

La, Asponse de diverges dispersions de carbure B la dtiformation cyclique a (5tk examink pour du fer tremp6 B foible teneur en oerbone, utilisant des essais de fatigue & amplitude de contrsinte constante, De plus, la r6ponse d’khantillons tremplts, dtiformb et vieillis, a tit.6 Btudii?t, pour d&erminer le tile de carbures dans la atabilisation de la microstructure de dislocation en fonction du ramollissement par fatigue. Las diff&emses dans lm disposition de dislocation a&s d&form&ion cyefique et d$formation unidireetionelle sont ration&s& en termes de dissipation de gradients de oontrainte- localis& sur 1s deformation invers6e. Les taux de ramvllissement par fatigue sent eonsid&& refl&er 1s grandeur des forces de retention exe&e sur la substruobure de dislocation par la dispersion des particules de la .seconde phase.

DIE STABILITjiT DER MICROSTRWKTUR IN DEM EISEN-KOHLENSTOFF- SYSTEM WAHREND ZYKLISCHER VERFORMUNG

Es ist die Reaktion versehiedener feiner H~~rne~ll-~v~~il~~n auf zyklis&e Yerformn~ fiir, dumb dbsohreekung gealterte, ni&@ Ei~~oble~~ffe unter Anwendung van E~~d~~pr~fun~n bei kvnetanter Spsnnungsamplitude untersucht worden. Ausserdem ist die Reaktion van abgemhreekt,eu, verfomten und gealterten Proben untersucht worden, urn die Rolle der Xerbide in der Stabilisieruug der Verschiebungsmikrostrulctur gegen Weichwerden dumb Ermiidnng r,u ermitteln. Die Unterschiede in Versohi?bun~ordmmg nai+ zyklifu?her Verformung und in einer Riohtung laufender Verfvrmung werden lm Smne der Dissipation 6rthoh festgelegter Sptmnungsgradienten auf riickl&nfiger Verformmg gedeutet. Es wird engenommen, dass die Grade des Erweichens dureh Err&dung die St&ke der suaek- haltenden K&f& wiedergeben, Welsh sich dureh die Dispersion der Teilcben aus der zwit;en phwe auf die Ve~h~eb~ngs~~~t~~ur auswirken.

INTRODUCTION

In many structural materials it is essentisl to produce not only a high yield strength in unidirectional drtformation bnt in ttddition resi&nee to failure during cyclic defo~&~ion. Previous investigations

have shown”*“’ that some materials strengthened by second phase particles suffer microstructural in- stabilities due to the localized removal of particles by cyclic deformation. Also, it is well established’3’ t.hat mat,&als deformed in unidire~ion~l deformation shows pronounced softening and microstructural change during subsequent cycIio deformation. In the past decade much effort has been given to pro. ducing combined work hardening and precipitation hardening mechanisms to increme the strength of engineering materi&. It is thus essential to under- stand both the role of secund phase particles during cyclic deformation and their effect on the stabilit,y of dislocation substructure produced during deformation prior to fatigue,

The objective of the present study was to delineate the role of iron earbide in controlling %he distribution

* Received July 30, 1971. t M&aster University, Hamilton, Ontario, Canada.

ACTA METALLURGICA, VOL. 20, APRIL 1972

of dislocations produced during cyclic strain and the stability of dislocation structures produced by prior deformation. The iron oarbon system w&s chosen in order to utilize a material in which the second phrase

particles are well characterized~~~ and easily varied in distribut,ion and also in order to study a model system which is of direct relevance to current. engineering materials.

EXPERIiSGD4TAL DETAILS

The study was performed on a low carbon iron of the composition shown in Table I, and of grain size approximately 50 ,u. In order to vary the distribution of second phase particles specimens were austenitized a& 85O”C, cooled to 725”C, W&S quenched snd aged at various ~ern~ra~ures between 25 and 240°C. In addition, some specimens were deformed by reduction of 8 per cent in diameter in a cold swaging operation after water quenching and prior to ageing and also after furnace cooling.

TABI;E 1

C Mn S N 0

0.018 0.037 0.013 0.003 0.072

628 ACT-4 METALLURGICA, VOL. 20, 1972

The microstructure of all specimens was studied by transmission electron microscopy both prior to and subsequent to the cyclic deformation in order to establish the detailed microstructural changes which occurred.

Cyclic deformation was performed using an M.T.S. closed loop eleetrohydraulic testing machine with a Woods metal grip system to minimize axial mis- alignment. The specimen was cooled during mounting and demounting to minimize structural changes due to heat Aow from the grip. The specimen gauge length was 0.5 in. in length and I/4 in. diameter and a11 specimens were carefully electropolished to reduce the influence of surface flaws. All tests were performed at a total plastic strain amplitude of AE, = AE = -&0.003 and at a frequency of 1 c/s. (In very long time tests the frequency was increased to 10 c/s.) In the interest of clarity, the present work is con- fined to descript,ions of changes in the bulk microst~~ture. The processes of local plastic instability and crack nucleation and growth are considered in a separate publication.

It is germane to note that in the plots of flow stress as a function of cyclic strain the irregularities due to the usual initial i~omogeneous yielding of the specimen are ignored. This aspect of cyclic hardening has been discussed in detail by Klesnil and Lukas.@j

RESULTS

For clarity the results are reported in three sections. The first describes the detailed microst~ctures observed prior to fatigue. The second cormerus the overall mechanical response to cyclic deformation and the final section describes the dislocation miero- structures produced by the cyclic deformation. It is of value in considering the microstructural features discussed in the third section to compare and contrast the observed dislocation microstructures with those generally observed in the unidirectional deformation of two phase materials.‘6-s’

In order to produce a wide range in particle distributions, specimens were aged at a variety of temperatures after quenching from 725%. After ageing at 250°C for 5 hr irregular distributions of dendritic particles were produced [Fig. I(a)}, Although no complete cross grating diffraction patterns were produced to identify the carbides from their morphology and habit plane they are consistent with the appearance of oementite. On ageing at 60°C for 6 days the precipitate distribution was extremely irregular due to the heterogeneous nucleation of

carbides on dislocations [see Fig. 1 (b)]. The tendency to precipitate on dislocations was found to increase with specimen size due to quenching strains and with the temperature from which the specimen was quenched due to the dislocations generated during the y to a transformation. The carbides on the dislocations are dendritic in form and appear to be cementite. A much more regular ~st~bution of 6ne particles was produced by ageing for 1 month at 25’C as illustrated by Fig. l(c). The (100) habit plane and the disc like morphology of t,he particles indicates that the precipitates are epsilon carbide discs approximately 400 _& in mean radius and 50-100 A in thickness. The dist~bution of carbides produced at 25°C was extremely uniform except, in the regions of grain boundaries and subboundaries where precipitate free regions of the order of 0.5 ,u in thickness were observed.

In addition to specimens which were aged directly after quenching, specimens were prepared which were swaged (by an 8 per oent reduction in diameter) between quenching and ageing. Specimens were aged at 60°C for 5 days and 240°C for 5 hr. Typical microstructures are shown in Figs. 2(a) and (b). It can be seen that the structure consists of a rough cell structure with regions of the order of 1 p in diameter separated by complex walls of dislocation. In the material aged at 60°C it is difficult to distinguish the precipitate particles on the cell walls but these are clearly visible in the material aged at 240°C. No quantitative assessment of the scale of the substructure formed by swaging and ageing was attemp- ted because of the complex shape of the cells which are formed in axisymetric deformation prooesses.@)

(b) Mechanical response

In order to provide a basis for comparing the behavior of the materials subjected to various ageing treatments and combined deformation and ageing procedures the flow stresses determined after various numbers of cycles have been normalized with respect to the initial flow stress for each specimen. In order to provide absolute values for the flow stress at any stage in the fatigue test reference can be made to Table 2 which summarizes the heat treatments, microstructural features and initial flow stresses of the material used in the present study.

The results shown in Fig. 3 indicate that the materials aged at lower tem~ratures had a higher resistance to fatigue softening than that aged at 240°C. In addition to the rate and magnitude of the fatigue softening the specimen aged at 240% had in general much lower fatigue lives than those aged at 25 or 60°C.

630 ACTA METALLURGICA, VOL. 20, 1972

.,

I

Fro. z(8). Microstructure after quenching, swageing and ageing at 60°C.

Fro. 2(b). Microstructure after quenching, swageing, snd sgeing at 24O’C.

Enhanced resistance to fatigne softening was observed in those materials quenched and swaged prior to ageing relative to materials swaged after slow cooling. It can be clearly seen from the data in Fig. 4 that the slowly cooled and swaged materials exhibit extensive softening during cyclic deformation, In contrast, the material swaged and aged at 60°C

shows little tendency to soften even after 108 cycles. Thus reference to Fig. 4 and Table 2 indicates that the material swaged to produce a ~sl~ation su~tNcture and aged at 60°C in order to stabilize the substructure shows an excellent combination of high yield strength and fatigue resistance.

The basic microstructural changes occurring as a

Series

QA 240

Pm-treatment

Quenched and aged 5 hr at 240°C

TABLE 2

Perticle size (rm)

Particle spacing Initial flow stress Fatique life (flm) (MN/m’) (cycles)

0.5 1.5 180 5 x 10’

QA 60 Quenched and aged 5 days at 60°C

0.5 0.4 230 105

S/A 25 Quenched and aged 1 month at 25’C

0.04 0.08 300 >2 x 106

F.C. Furnace cooled 150 4 x IO”

F.C. and S Furnace cooled and 200 5 x 10’ swaged 8 y0 R. Of. D.

Q and A 240 Quenched and swaged -1.0 220 5 x 10’ 8 % R. Of. D. aged 5 hr at 24O’C

Q and A 60 Quenched and swaged -0.1 230 2 x 10s 8 % R. Of. D. aged 5 days at 60°C

BROWN% AND EMBURY: STABILITY DURING CYCLIC DEFORMATION 631

I I ,

K1- 3z -

wnwz” OF c*cl.16*

Fro. 3. Cyclic softening curves for quenched end aged specimena.

result of cyclic deformation are described in the next

section.

(6) Yicrostructures after fatigue

In the specimens aged at 24O’C which contained coarse distribution of cementite it was observed after some 5000 cycles that dislocations were accumulated in the vicinity of the particles. No local rotations were produced in the vicinity of the particles which is in marked contrast to the behavior normally observed in unidirectional deformation. After lo” cycles a cell structure was produced which was similar in scale to that produced in a specimen containing no second phase particbs, as illustrated in Figs. .#(a) and (b).

In the material aged at 60°C the complex arrays of precipitates and grown in dislocations served as sites for the accumulation of large densities of loops and dipoles during fatigue as shown in Fii. 6, The lack of misorientations again indicates that the accumulated dislocations must be essentially equal numbers of dislocations of opposite sign.

The specimens aged at 25°C and containing fine dispersions of particles showed little change after fatigue. Some dipoles were accumulated in the vicinity of the particles and a very simple cell structure was produced as illustrated by Fig. 7. The cell walls

were simple in form and the soale of the structure appears to be independent of the particle distribution and was observed only after about 108 cy&s.

None of the structures taken from the bulk of the quenched and aged materials showed evidence of particle dissolution. However, some regions taken from close to the crack tip and to the specimen surface did show evidence of highly localized plastic flow, although the observed structures did not indicate unambiguously that localized flow, occurred as a consequence of precipitate dissolution. This aspect of the microstructure is considered in detail in the publication concerning the stability of plastic flow.

Fm. 4. Cyclic eoftening eur~es for quenched and swaged specxxueue.

The material which was swaged and aged at 240°C eventually showed some fatigue softening and the irregular cell structure illustrated in Fig. 2(b) was converted to a more regular arrangement of dislocation walls as shown in Fig. 8. At higher magnifica- tions large numbers of dipoles could be seen associated with the regular dislocation walls. Similar microstructural changes were observed in the material swaged and aged at 60°C but only after a very much larger number of cycles.

DEKX.JSSIOI’3

The salient features of the results reported in the previous section may be summarized as follows:

(a) Second phase particles do not promote rapid hardening in cyclic deformation and the resultant dislocation subst~otures differ markedly from those produced in uniaxial deformation.

(b) The scale of the dispersion of second phase particles is important in determining the form of the accumulated dislocation distribution. Coarse dispersions of particles appear to promote a cell structure but fine dispersions exert little influence on the form of the diilocation substructure.

(c) Combinations of deformation and ageing promote increased fatigue resistance due to the stability of the dislocation substructure produced by the combined deformation and ageing procedure.

In order to rationalize these features, let us consider firstly the basic geometric requirements of a fatigue test. In cyclia deformation the mean strain is zero and thus second phase particles cannot impose any additional prooess of dislocation accumulation on the matrix as they can in uni~~ctional deformation.~~l~ Turbulent flow may occur locally in the vicinity of the part&lea on the first tensile half stroke but the results of this localized flow will be dissipated in subsequent cycles.02) Hence, dislocations will ac. cumulate at the particles in the form of dipoles with no


FIG. 6(a). Accumulation of dislocations and formation of a sub-boundary ctfter 6000 cycles, in a specimen quenched

and aged at 240°C.

7. ._. _“- -

Fig. 5(b). Cell structure formed in the same materiel after 10” cycles

Fro. 6. Accumulation of dislocations at 6.6 X IO6 cycles in materiel aged at f3O’C.

SItOTVN~ AND EMBURY: BTABILITY ,D,URXNC CYCLIC DEFORMATIOX 633


resultant lattice rotation as observed in Fig. 6. A% the stress field of dipole% falls off more rapidly withdietance than that of single dislocations dipole% are fairly in- effectivein hardening the ory%tal. Thus, in cyclic deformation the hard second phase particles do not promote rapid work hardening as in unidirectional deformation and no local lattice rotation% are produced at the particles in cyclic deformation in contrast with those observed in tensile deformation of two phase materials.

Let u% now consider some basic aspects of the influence of second phase particrles on the accumulation and retention of dislocation subStrUCtUre. When particles are small they can be effectively by-pas%ed by cro%% slip, and thus the particles will exert little influence on the subsequent accumulation of dislocations and the cell structure which is formed will be independent of the dispersion of the particles as in Fig. 7. Similar results concerning the role of small particles have been observed in un~ctional deformation at large plastic strain%.(fs) A simple dimensional argument can be advanced in regard to the scale of the substructure by considering a dislocation held up at a line of particles and unable to escape by cro%% slip or climb. The dislocation attached to the particles will be subject to forces from the presence of other d&locations at a distance t which are of the order of pbe/2nr. However, due to the presence of the particles the dislocation is effectively subject to a restraining force due to the particle which is equal in magnitude to the Orowan stress, For particles spaced a distance il apart the restraining force is ,ub2/2A. Thus, it is expected that when cells form in two phase materials with large particles which are not by-passed by oroas slip, the cell dimen- sions will be of the order of r and will thus be of the order of the particle spacing a% is generally observed. (These simple conditions will not apply if the cell walls consist of groups of like dislocations or dipoles.)

One of the most salient features of the present study is the role of small precipitates in stabilizing existing dislocation %u~tNcturea against cyclic softening (Fig. 4). In general, cyclic softening involves the formation of large numbers of dipoles from the dislocations accumulated during the prestrain, the pro- duction of a symmetrical hysteresis loop and a marked decrease in the local lattice rotations produd by the prestrain.(i’) In order for these prooessea to occur the dislocation% accumulated during preatrain must interact with the dislocations moved back and forth during cyclic deformation to form a simple dislocation structure as shown by comparison of

Figs. 2 and 8. The dominant forces causing softening must be the elastic interaction of dislocations and thus the ~sl~tio~ accumula~ during prestrain must experience a foroe of the order of pbzj2w due to a moving dislocation at distance r. In order to resist interaction with such dislocations the acoumu- lated dislocation% must be pinned by a resistive stress greater than $%/2nr. This can easily arise due to the formation of a solute atmosphere or a series of fine precipitates along the dislocation line (as in the case of ageing at 60°C after prestrain). The resistive stress on the accumulated dislocation will be of the order of (l/b)(a~~~~~~x) where V,, is the interaction energy between the atmosphere and the ~slo~tion. Typic- ally, suoh resistive stresses will be of the order of 5-20 kg/mm2 even for dilute atmospheres or arrays of small particles along the dislocation line and thus the accumulated dislocation structure will exhibit enhanced resistance to fatigue softening as seen in the data presented in Fig. 4.

The above discussion illustrates the essential differences between the behavior of two phase material% in fatigue and unidirectional deformation and in addition provides a simple rationalization for the mechanism of thermal mechanical processing in improving the stability and fatigue resistance of materials. Similar arguments may be applicable to more complex commercial alloys such a% materials subject to dynamic strain ageing and to ausformed alloy steels. The basic argument concerning the stability of dislocation substructure can al80 be extended to include the climb forces of dialoca- tions to illustrate some basic similarities between the mechanical stability of dislocation substruo- tures and their resistance to thermal recovery.

The present study has been concerned with the bulk behavior of two phase iron-iron carbide systems in regard to both mechanical behavior and microstructural changes. No local particle dissolution wa% observed in the bulk and the role of localized plastio deformation and the associated microstruotural instability are more relevant to crack propagation and is discussed in a separate publication.

ACKNOWLEDGMENTS

The authors wish to thank the American Iron and Steel Institute for their generous research support and the Nations Be%eareh Council of Canada for the provision of a research studentship (M. J. B.). One of the authors (J. D. E.) would like to express his thanks to Battelle’s Columbus Laboratories for the pro&ion of a Battelle Institute Fellowship during which the present study wa% completed.

BROWPU’# AND EMBURY: STABILITY DURING CYCLIC DEFORMATIOX 635

REFERENCES

1. P. J. E. FORSYTH and C. A, STUBBINGTON, J. Inst. Metale 90, 347 (1961).

2. J. T. MCGRATH and W. J. BRATIXA, Acta Met. 15, 329

3. !‘?&.TNER a.nd C LAI~D Acta Met. 15 1621 (1967) 4. G. c’. LESLIE, R. 113. F&HER akd N. SEX, .&a Met. 7,63;

(1959). 5. M. KLESNIL and P. LUKAS, J. Iron Steel Inst. 205, 546

(1967). 6. A. KELLY and R. B. NICHOLSON, Prog. MetaE Pkyls. 10,

(1963).

7. K. C. RUSSELL and M. F. ASHB~, Acta Met. 18, 891 (1970).

8. P. B. HIRSCH and F. J. H~TNPHRETS, Physics of Strength and PEasticity, edit.ed by A. S. AEGOX-, p. 189. M.I.T. Press.

9. G. LANDLORD and M. COTTEN, Trans. Am. Sot. Metals 62, 623 (1969).

10. M. F. ASHBY, Phil. Mag. 15, 1157 (1966). 11. W. M. STOBBS, D. F. WATT and L. M. BRO~\‘IC, Phil. Xag.

2$,1169 (1971). 12. J. D. EMBURY, A. S. KEH and R. 31. FISHER, Trajw.

Am. Inst. Min. Engrs 295,1252 (1966). 13. J. D. EMBVRY and D. HTZG~ARD, to bp published.

the stability of microstructure in the iron-carbon system during cyclic deformation

Documents